Method of aging iron-base austenitic alloys



T. w. EICHELBERGER 2,879,194

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METHOD OF AGING IRON-BASE AUSTENITIC ALLOYS Filed July 12, 1957 3Sheets-Sheet 2 Fig.2.

- I I 1 I Percent Titanium Notch Creep- Rupture Tests of AusieniticAlldys March24, 1959 T. w. EICHELBERGER 2,879,194

METHOD OF AGING IRON-BASE AUSTENITIC, ALLOYS Filed July 12, 1957 5Sheets-Sheet 3 Aeme TREATMENT 2 4 7 0 O Plain-Bar Creep-Rupture Numbersv are Percent Elongation for Plain Bar Tests ipoo N6tched 60,000 psi LRupture Time in Hours m an 6 o o o o. l

I i l I L L4 L6 L8 2.0 2.2 2.4 216 2.8

Percent Titanium I200 F Creep-Rupture Tests on Austenitic Alloys withVarying Titanium Contents Fig.3.

United States Patent METHOD OF AGING IRON-BASE AUSTENITIC ALLOYS ThomasW. Eichelberger, Pittsburgh, Pa., assignor to Westinghouse ElectricCorporation, East Pittsburgh, Pa., a corporation of PennsylvaniaApplication July 12, 1957, Serial No. 671,598

7 Claims. (Cl. 148-2154) This invention relates to a heat treatmentprocess for the hardening of precipitation hardening alloys.

At the present time high temperature age-hardenable alloys are employedextensively for applications in gas turbines and jet engines. Inattaining the desired physical properties such as creep-rupture strengthand duetility, members of the wrought alloys are subjected to selectedheat treatment procedures. Such previously employed heat treatmentprocedures usually have comprised the heating of the member initially toa solution treatment temperature of from about 1700 F. to 2200 F., andthen the members have been subjected to either a double aging treatmentor a single aging treatment. The known commercial double aging heattreatments have resulted in an excess of age-hardening precipitateforming both at grain boundaries and at imperfections in the crystalstructure. Consequently, the alloys have not been completelysatisfactory because the excessive grain boundary precipitate hasresulted in low ductility on 2,879,194 Patented Mar. 24, 1959 quiringsuch lengthy periods of time add greatly to the cost of the wroughtmembers and are not commercially acceptable.

The object of the present invention is to provide a 5 multiple stageaging heat treatment procedure for precipitation-hardenable hightemperature alloys which may be carried out in a relatively short totalperiod of time and which will develop optimum creep-rupture propertiesfor the given alloy.

A further object of the invention is to provide an aging treatment forhigh temperature austenitic alloys containing .a precipitation-hardeningcomponent wherein the alloys after solution treatment are heat treatedinitially in a temperature range of from 1100 F. to 1250 F. for a periodof time of the order of from 2 to 50 hours and then slowly raised to ahigher temperature of from 1300 F. to 1450 F. and maintained in thislatter temperature range to attain optimum creep-rupture prop:

erties.

creep-rupture tests. To avoid difiiculties, the amount ofprecipitation-hardening component in the alloy is critically restricted.g

The previously employed single aging heat treatment has not beenentirely satisfactory for a number of reasons. A primary drawback hasbeen the fact that such single heat treatments have required excessivelylong aging times to secure the best creep-rupture properties obtainabletherefrom. Atleast 100 hours and as much as 1000 hours in the agingfurnace is required to secure good creep-rupture properties for thegiven alloy.

The maximum hardness and the optimum creep-mp A stillfurther object ofthe invention is to provide a twostage aging treatment of asolution-treated wrought member of a high temperatureprecipitation-hardening austenitic alloy wherein the first stepcomprises aging the alloy in the range of from 1100 F. to 1250" F. tocause nucleation of the precipitation-hardening component whereby toincrease the hardness of the member and then slowly increasing thetemperature of the member to from 1300 F. to 1450 F. and maintaining themember in this last temperature range to cause the full hardness of themember to be developed with very little grain boundary precipitate.

Other objects of the invention will in part be obvious and will in partappear hereinafter.

For a better understanding of the nature and objects of the invention,reference should be had to the following detailed description anddrawing, in which:

Figure 1 is a graph plotting hardness against the total aging time of analloy processed in accordance with the present invention;

Fig. 2 is a graph plotting notch bar creep-rupture time against thetitanium content of a given precipitationhardened alloy following thepractice of the present ture properties for a given .alloy being given'a single aging treatment can be secured only at the lower agingtemperatures applied for very long periods of time; For example, when agiven titanium hardened austenitic alloy was aged 400 hours at 1200 F.,it had a creep-rupture time of 108 hours at 60,00 p.s.i. at 1200 F., andan elongation of 17.2%. Other specimens of the same alloy aged for 20hours at 1300" F., when tested at 60,000 p.s.i. loads at 1200 F. had acreep-rupture time of only 19 to 35 hours with the elongation being from15% to 17%. This example emphasizes the long aging times re-' quired toproduce an above average aged product by a single aging treatment.

It will be appreciated that single aging treatments reinvention andfollowing the practice of the prior art; and Fig. 3 is a graph plottingrupture time in hours against the 1200? F. creep-rupture properties of aprecipitationhardened alloy in which the titanium hardener content isvaried.

The present invention particularly relates to high temperatureaustenitic alloys in which titanium alone or in combination with othercomponents functions as the precipitationrhardening component. Morespecifically,,these alloys will contain'from 1.3% to 3.5% by weight oftitanium. The alloys may, and usually will, also contain molybdenum and/or tungsten, and aluminum which exhibit some precipitation-hardeningcharacteristics, though not to as marked a degree as that of thetitanium.

Broadly, the austenitic alloys to which the present invention isparticularly adapted are composed of from 10% to 35% by weight ofnickel, from 7% to 23% by.

Weight of chromium, a total of up to 6% by weight of at least one metalfrom the group consisting of tungsten and molybdenum, up to 2.5% byweight of manganese, between 0.3% and 1.5% by weight of silicon, up to0.5 by weight of aluminum, up to 0.5% by weight of vanadium, up to 0.5%by weight of boron, from 1.3% to 3.5 by weight of soluble titanium, lessthan 0.15% of carbon, and the balance being iron except for incidentalimpurities such as phosphorous, sulfur, nitrogen, oxygen and the like. Adesirable range of alloy compositions for processing in accordance withthe invention comprises from 22% to 28% nickel, from 12% to 18%chromium, a total of from 1% to 3% of at least one metal from the groupconsisting of tungsten and molybdenum, from 0.5% to 1.5% manganese, from0.1% to 0.35% vanadium, less than 0.10% carbon, from 0.8% to, 1%silicon, up to 0.5% aluminum, from 1.5% to 2.8% soluble titanium, from0.03% 'to 0.15% boron, and the balance being iron except for incidentalimpurities. Outstanding properties have been obtained in heat treatingalloys containing from 22% to 28% by weight of nickel, from 12% to 18%by weight of chromium, from 1% to 3.5% by weight of molybdenum, from0.03% to 0.15% by weight of boron, from 0.1% to 0.50% by weight ofvanadium, less than 0.08% carbon, from 1% to 1.5% by weight ofmanganese, from 0.8% to 1% of silicon and up to 0.4% by weight ofaluminum, the balance being iron except for incidental impurities. Ifthe alloys are vacuum melted lesser amounts of manganese, silicon andaluminum may be employed to a'ccomplish the usual scavenging functionsthereof. Soluble titanium may be determined, for practical purposes, bysubtracting from the total titanium content an amount equal to 4 timesthe carbon content of'the alloy.

The alloy may be prepared by induction melting either in the open air orunder a protective atmosphere, or unu'er'v'ac'uum, or by are melting inan evacuated vessel or with'a preteaive inert gas such as helium orargon. Ing'ots or castings may be prepared from the alloy. The ingots orcastings may be subjected to forging, hot rolling, contour die, forgingor other appropriate hot working procedures. The resulting wroughtmembers which may be in the form of discs, bar stock, bolts, sheets andthe like, must be subjected to an aging heat treatment in order todevelop the physical properties 'to the'desired extent.

In accordance with the present invention, the wrought members of theprecipitation-hardenable austenitic alloy are initially solution treatedat a temperature above 1700 F. In practice the solution treatment may becarried out-at a temperature from 17.00 F. to 2200 F. in a period oftime of from $6 to 8 hours. After soluand preferably from 4 to 20 hours.The temperature need not be constant in this range. During this firstaging treatment it has been discovered that very fine, evenlydistributed particles of the precipitation-hardening components arenucleated. There is no excessive segregation of the hardening componentsnear the grain boundaries or other imperfections. The extent of thisinitial nucleation precipitation can be measured by the increase inhardness of the alloy member. The hardness will usually increase from 30to 125 points as determined by the Diamond Pyramid Hardness Test,abbreviated as DPH or VHN (Vickers Hardness Number). Ordinarily, thefirst aging treatment has given excellent results when the hardness hasincreased from 50 to 100 DPH.

After the nucleation of the precipitation-hardening components has takenplace to the desired extent, the alloy member is then slowly increasedin temperature to a second aging temperature of from 1300 to 1450 F. Inorder 'to prevent redi's's'olutibn of some of the fine hardeningprecipitate, the temperature should be raised relatively uniformly andfor a prolonged period of time of from 2 to 32 hours to the desiredsecond aging temperature.

Once the desired second aging temperature has been reached, the memberis held or maintained at this second aging temperature until it hasattained the desired optimum hard ness. The rate of increase of hardnessof the alloy is much more rapid than during the first aging temperature.The hardness should increase at least 25 DPH points. The ordinarilydesired optimum hardness of the members after the second aging heattreatment will be at least 225 DPH for an alloy containing 1.3% solubletitanium. and exhibit progressively higher hardness with increasingsoluble titanium content to 330 to 360 DPH for an alloy containing 2.4%soluble titanium.

After the alloy member has been subjected to the second "agingtemperature, it may then be cooled to room temperature or otherwise 'putinto condition for use. 7

The aging treatment of this invention is designated as ge'nuage heattreatment.

In some cases it may be desirable to secure a further increase ofhardness for the given alloy. To accompli'sh this, it has beendiscovered that by reducing the temperature of the alloy member from thesecond aging temperature, that, is from 1300 F. to 1450 F., to atemperature of about 1200" F. and holding the member for a period oftime of at least 2 hours at this third temperatur'e, a marked increasein hardness will occur.

Heats of the several alloys shown in the attached Table I were preparedby melting the components in an induction. furnace, in, air.

TABLE I, Ghem'ical compositionef alle s- Ni Cr Mn s1 0 Al v Fe 2413 14.21.38 .76 .036 11 .068 Balance. 24.8 4.1 1.38 .82 .032 .-12 .078 Do.524.3. 14.1 1. 41 87 -.032 .1-1. .078 -Do. 25.2 14.14 1.28 .74 .024 .10Do. 25.5. 14.2 1.28 .74 .02 .19 Do. 25.2 14.5 1.17 .88 038 .06 D0,, 26.413.7 1.60 .71 .04 .15 .08 Do.

Ingots'Were-"prepared from each of. the alloys of Table I. V The. ingotsw'eresoake'd for four hours at 2100 F., and then subjected to 8 seriesof hot rolling stages with intermediate reheating, to produce /B-lIlCl'lround bar. Specimens were machined from the round bar and sub,- jectedto a variety of heat treatments. Table II indicates the variety of heattreatments of this invention applied to the alloy of each specimen.

of heatNo. 7040, and the resulting hardness TABLE n Aging cycle tests 11 hour at 1950 F.

[Hardness in DPH; time in hours.]

Hardness Heating Hardness Aging Time at 1,300 F. Specimen as Solu- AgingHardness Time After Number tion Time at Alter from Reaching Treated1,200 F. 1,200 F. 1,200 to 1,300 F. 2 4 8 12 16 20 129 4 197 2 195 219230 274 12s 4 192 4 202 277 130 4 192 3 209 238 242 278 126 4 192 16216' 242 243 286 130 3 198 2 191 27s 128 v s 198 4 196 276 130 3 200 8206 250 276 130 3 196 16 217 247 274 129 16 211 2 196 270 130 16 210 4203 276 12s 16 209 3 208 240 246 284 130 16 209 16 229 224 241 272 12732 221 2 206 272 12s 32 21s 4 211 282 130 32 223 s 216 273 131 32 220 16229 274 129 64 236 2 222 132 64 234 4 224 130 64 234 s 226 127 64 239 16248 268 Referring to Flg. 1 of the drawlngs, the curves there TABLE IIIshown graphically illustrate the data of Table H, 1n terms Creep testsof heat N0. 7040 at 12003 R All specimens of the change in hardness ofthe alloy of heat No. 7040 which contains 1.69% total titanium as themajor precipitation-hardening ingredient and 3.0% molybdenum whichcooperates therewith. The curves on'Fig. 1 show the change in hardnessas a function of time in hours for the wrought members (a) during thefirst aging treatment, (b) during the period of time when thetemperature is being changed from 1200 to 1300 F., and (c) finally whileaging at 1300 F. Thus, it will be noted that the initial hardness of thealloy is approximately 130 D.P.H. before any aging treatment. Afterapproximately 4 hours at 1200 F., the hardness has increased to about195 D.P.H.an increase of approximately 65 points D.P.H. A slow increasein hardness occurswith further aging after 4 hours at 1200 F. so thateven after 64 hours aging time it has only reached 235 D.P.H. It will benoted that as the period of time consumed in raising the temperature ofthe alloy from 1200 to 1300 F. varies from 2 to 16 hours, there is aninitial slight drop in hardness, which is believed to be due to some ofthe hardening precipitate redissolving, soon, however, an increase inhardness is observed. A rather high rate of increase in hardness occurswhile the alloy is being heldv at 1300 F.

It will be noted further that when the first aging heat treatment at1200 F. exceeds 8 hours, there is a definite.

drop in hardness if the temperature is increased from 1200 to 1300? F.in from 1 to 8 hours. However, there is no decrease in hardness if aperiod of time of the order of 8 to 16 hours is used to increase thetemperature from 1200 to 1300 F.

Prlonged aging times employing the three stages of heat treatment arenot required to attain optimum hardness. hours, the hardness of thealloy is approximately 275 D.P.H., while for a total aging time of 40hours, the maximum hardness is approximately 285 D.P.H. and for 84 hoursit attains only 265 D.P.H.

Since the ultimate hardness of the alloy, regardless of the length ofthe total heat treatment time, is between about 265 and 285 D.P.H., hereis no need to employ any heat treatment schedule that exceeds 40 hours.Figure 1 of the drawing therefore emphasizes the fact that the processof the present invention produces outstanding hardness for a relativelyshort total aging time of between 20 and 40 hours. 7 Creep-rupturetestsof the heat treated alloy No. 704 are set forth in the followingTable III.

Thus, for a total aging time of approximately 26 solution treated 1 hourat 1950 F.

PLAIN-BAR SPECIMENS-AT 60,000 P.S.I.

Rupture Minimum Aging Hard- Creep Specimen Treatness as Rate, Numberment Aged Time, Elonga- Percent DPH Hours tion, Hour I Percent 289 6910.0 OObfi 230 1 9 9.8 6'04 283 181 10. 3 0. 017 272 141 15.3 0.02

NOTCH-BAR SPECIMENS-AT 70,000 P.S.I.

Aging Hardness Time, Specimen Number Treatment as Aged Hours DPH Theaging treatment applied to the specimens :is im dicated in the attachedTable- IV. It will be observed that aging treatments A, B and C arethose of the present invention, while treatment D is a standardprior artheat treatment in wide use. Aging treatments A, B and C resulted in twoto three times'the rupture time for plainbars as obtained with treatmentD. In notch-bars, aging treatments A, B and C result in 5 times or morerupture time as from treatment D, while the elongationisas good orbetter in A, B and C as compared to aging treatment D.

TABLE IV Aging Treatment A 4 hours at 1,200 F., temperature increased to1,265? F. in 1% hours, held at 1,265" F. for 2 hours, raised to 1,300 F.in 10 minutes, held at 1,300 F. for 20 hours and air cooled. t B 8 hoursat 1,200 F., temperature increased to 1,300 F. in 8 hours, held at 1,800F.1or 20 hours and air cooled. O 16 hours at 1,200 F., heated to 1,300F. in 16 hours,.held

at 1,300 F. for 12 hours and air cooled. 20 hours at 1,300 F. cooled to1,200 F. in 4 hours, held at 1,200 F. for 20 hours and air cooled.

7 The eflz'cct of the aging treatment of the present invention ascompared to other widely used aging treatments for various specimens ofsome of the alloys of Table I is set forth in the attached Table V.

8 of the present invention. Further, it will be noted that alloys may beprocessed in accordance with the present invention with much highertitanium contents then would be considered usable when aged by priorprocedures.

TABLE V The efiect of various aging treatments on the creep propertiesof some ausfenitic alloys at 1200 F. and

60,000 p.s.l-

Hardness Transition Time to Given Strain in Hours Rupture 7 Red. MinimumAging in Creep Specimen Number Treat- After Elonga- I Elonge- Area,Rate, Perment AsAged Test- Time, tion 03% 0.5% 1.0% 211% Time, tion,Percent/Hour DPH ing Bours Percent Hours Percent cent DPH Heat treatment1 is a standard heat treatment widely used at the present time. Itcomprises aging for 20 hours at 1350 F., cooling to 1200 F. in 4 hoursand holding at 1200" F. for 20 hours followed by air cooling. Heattreatment 1* is the same as heat treatment 1 except the time to coolfrom 1350" F. was carried out in 3 hours. Heat treatment 5, anotherwidely used commercial heat treatment for these alloys, comprises agingthe alloy for 16 hours at 1325 F. Heat treatment 4 constitutes a'heattreatment in accordance with the present invention wherein the alloy isheated for 16 hoursat 1200" F., heated to 1300 F. over a period of 16hours, then held at 1300 F. for 12 hours and air cooled. Heat treatment4* comprises heating 16 hours at 1200 F., then increasing thetemperature over a period of 8 hours Thus, for alloys having a 2.3%titanium content, the

a notch-bar creep-rupture time is of the order of 2 hours when aged byeither the single or double aging treatments of the prior art. Using theaging treatment of the present invention, a 2.3% titanium alloy exhibitsa creeprupture time of approximately 280 hours. The curves on Fig. 2therefore emphasize the outstanding improvement inthe alloy propertiesobtainable by the practice of the present invention.

TABLE VI The efi ect of various aging treatments on the notch-barcreep-rupture properties of some austenitic type alloys at 1200 F.

to 1250" F. and holding at 1250 F. for 16 hours and air cooling. 7 AgingHardness Rupture With respect to alloys having low soluble titamumSpecimen Number Treat- Load, Time, content of below 1.4%, it ispreferred to modify slightly gg Hmus the second aging temperature to arange of from 1250 DPH DPH F. to 1350 F. This corresponds to heattreatment 4*. The first. aging temperature. range for such low titanium1 297 299 70,000 180 alloy should be preferably from 1100 F. to 1200" F.g 5g? :838 38 Best aged alloys having soluble titanium contents of 4 277295 70,000 403 1.3% to 1.4% have been obtained with such modified 3;;Q883 3g e fls 5 295 329 so, 000 261 It will be observed from Table IIthat the first four 55 $2 22% 53 833 f? alloys exhibited much superiorcreep-rupture times when 1 333 349 60.000 23 heat treated in accordancewith the present invention than Z 32; 381888 5.2 when the same alloy washeat treated by other known 5 320 347 ,03 153 processes. The ductilityas measured by the elongation ,3 18 3; was favorably comparable to thatof the members with 1 24 8 65.000 199 in 'mp time's it; 332 3%;338 itNotch-bar creep-rupture tests established that the alloys aged by theprocess of. the present invention were less notch sensitive than thesame alloys processed by other commercially used aging treatments. Theloads inmany cases had to be increased to 70,000 p.s.i. to securerupfine iii 'rea'sfinabl'e periods or time.

Referring; to Fig. 2 of the drawings, there are plotted curves based ondata obtained from a considerable number ofnotch-bar. creep-rupturetests of similar alloys with varying total titanium contact. It will benoted that the notch-bar'creep-rupture times using the single and doubleaginghe'at. treatments of the prior art are considerably inferior to thecree -rupture times, for similar alloys with the same titanium content,obtained by the aging process A summary of the various physicalproperties obtain able by the practice of the present invention onvarious precipitatiomhardenable alloys as compared to those obtainableby the best prior art practices, for varying titaniu'm content ispresented the curves of Fig. 3. Aging treatment 2 of Fig. 3 comprises 20hours at 1300 F., cool to 1200" F. in 4 hours, hold at 1200 F. for 20hours and. air cool; it has been widely used. Aging treatment 4 is thatof the present invention and is that set forth above in Table V. Theaging treatments of Fig. 3 were-applied to alloys corresponding tothosein Table I.

. It will be observed that aging treatment 4 results in memberswhosecreep-rupture times at stresses of 70,000 p.s.i. exceeds by asubstantial amount that of alloy members aged by treatment 2 but loadedto only 60,000 psi. It will be apparent that even at these higherstrengths, the ductility of the alloy members is superior in generalwhen aging treatment 4 of this invention is employed as contrasted tothe prior art heat treatment. I i

The following alloy was melted and cast into an ingot: nickel 26.1%;chromium 13.1%; molybdenum 3.13%; carbon 0.018%; silicon 1.01%;manganese 1.33%; titanium 2.4%; boron 0.12%; and 52.4% iron, balanceincidental impurities. The ingot was forged at 2000'F. to bars whichwere hot rolled to 78 diameter stock. The bars were solution treated onehour at 1950 F., and quenched in oil. The grain size of the solutiontreated alloy was 4 to 5 A.S.T.M. Part of the bars were aged by heatinghours at 1350 F., cool to 1200 F. in 5 hours, hold 20 hours at 1200 F.,and air coolthis is indicated as treatment X. Other bars were aged byheating at 1200 F. for 16 hours, slowly increasing the temperature to1300 F. over 16 hours, holding 12 hours at 1300" F. and finally aircooling. This latter is aging treatment Y. Creep-rupture tests at 1200F. at 70,000

p.s.i. load resulted in the following results:

Notched Hours to Elongation, ar, Aging Treatment Rupture Percent Hoursto Rupture It will be understood that the above description and drawingare illustrative and not in limitation of the invention.

I claim as my invention:

1. In the process of producing a precipitation hardened member from aniron base austenitic alloy containing from 1.3% to 3.5% by weight ofsoluble titanium as a precipitation hardening component, the stepscomprising (1) solution treating the alloy member at a temperature above1700 F., (2) rapidly cooling the alloy member to a temperature below1250 F., (3) reheating the cooled member to a first aging temperature offrom 1100 F. to 1250 F. and preferably between 1100" F. and 1200 F. foralloys having less than 1.4% of soluble titanium, and maintaining themember in this temperature range for a period of time of from 2 to 50hours sutficient to cause nucleation of the precipitation hardeningcomponents whereby the hardness of the alloy increases from to 125points by the diamond pyramid hardness test, (4) then heating the alloymember over a period of from 2 to 32 hours to increase its temperatureslowly to a temperature of from about 1300 F. to 1450 F., and preferablyfrom 1250 F. to 1350" F. for alloys having less than 1.4% of solubletitanium, and maintaining the alloy member in this range until itincreases in hardness at least 25 points more than the hardness afterthe first aging temperature.

2. The process of claim 1 wherein after the heating at from 1300 F. to1450" F., the temperature of the member is reduced to about 1200 F. andmaintained at this temperature for a period to produce a furtherincrease in hardness.

3. In the process of producing a precipitation hardened member from anaustenitic alloy composed of from 10% to 35% nickel, from 7% to 23% ofchromium, a total of up to 6% of at least one metal from the groupconsisting of tungsten and molybdenum, up to 2.5% manganese, between0.3% and 1.5% silicon, up to 0.5% aluminum, up to 0.5% vanadium, lessthan 0.15% carbon, up to 0.5% boron, from 1.3% to 3.5% soluble titanium,

10 the balance being iron except for incidental impurities, the stepscomprising (1) solution treating the alloy member at a temperature above1700 F., (2) rapidly cooling the alloy member to a temperature below1250' F., (3) bringing the member to a first aging temperature of from1100 F. to 1250 F. and maintaining the alloy member at this temperaturefor a period of from 2 to 50 hours to cause nucleation of theprecipitation hardening components whereby the hardness of the alloyincreases from 30 to points by the DPH test, and (4) thereafter heatingthe alloy member over a period from 2 to 32 hours to increase itstemperature slow- 1y to a second aging temperature of from 1300 F. to1450 F. and maintaining the member at this second temperature to causethe hardness to increase at a more rapid rate than when at the firstaging temperature to a value proportional to at least from 225 DPH for a1.3% soluble titanium alloy to 330 DPH for a 2.4 3% soluble titaniumalloy.

4. In the process of producing a precipitation hardened member from anaustenitic alloy composed of from 10% to 35% nickel, from 7% to 23% ofchromium, a total of up to 6% of at least one metal from the groupconsisting of tungsten and molybdenum, up to 2.5% manganese, between0.3% to 1.5% silicon, up to 0.5% aluminum, up to 0.5 vanadium, less than0.15% carbon, up to 0.5% boron, from 1.3% to 3.5% soluble titanium, thebalance being iron except for incidental impurities, the stepscomprising (1) solution treating the alloy member at a temperature above1700 F., (2) rapidly cooling the alloy member to a temperature be low1250 F., (3) bringing the member to a first aging temperature of from1100 F. to 1250 F. and maintaining the alloy member at this temperaturefor a period of from 2 to 50 hours to cause nucleation of theprecipitation hardening components whereby the hardness of the alloyincreases from 30 to 125 points by the DPH test, (4) thereafter heatingthe alloy member over a period of from 2 to 32 hours to increase itstemperature slowly to a second aging temperature of from 1300 F. to 1450F. and maintaining the member at this second temperature to cause thehardness to increase at a more rapid rate than when at the first agingtemperature to a value proportional to at least from 225 DPH for a 1.3%soluble titanium alloy to 330 DPH for a 2.4% soluble titanium alloy, and(5) thereafter reducing the temperature of the alloy member to about1200 F. for a period of time of at least 2 hours to produce a furtherincrease in hardness of the alloy.

5. In the process of producing a precipitation hardened member from anaustenitic alloy composed of from 10% to 35% nickel, from 7% to 23% ofchromium, a total of up to 6% of at least one metal from the groupconsisting of tungsten and molybdenum, up to 2.5% manganese, between0.3% and 1.5 silicon, up to 0.5% aluminum, up to 0.5% vanadium, lessthan 0.15 carbon, up to 0.5% boron, from 1.3% to 3.5% soluble titanium,the balance being iron except for incidental impurities, the stepscomprising (1) solution treating the alloy member at a temperature above1700 F., (2) rapidly cooling the alloy member to a temperature below1250 F., (3) bringing the member to a first aging temperature of from1100 F. to 1250 F. and maintaining the alloy member at this temperaturefor a period of from 2 to 50 hours to cause nucleation of theprecipitation hardening components whereby the hardness of the alloyincreases from 30 to 125 points by the DPH test, and (4) thereafterheating the alloy member over a period of from 2 to 32 hours to a secondaging temperature of from 1300 F. to 1450 F. and maintaining the memberat this second temperature to cause i the hardness to increase at a morerapid rate than when at the first aging temperature to increase itstemperature slowly to a value proportional to at least from 225 11DPH'for a 1.3% soluble titanium alloy to SSODPH for 2121492; solubletitanium alloy.

6. In theprocess' of producing a precipitation hardened member' froman'austenitic alloy composed of from 22% to 28% of'nickel, from 12% to18% chromiunna total offrom 1% to 3% of at least one metal selectedtromthe group consisting of tungsten and'molybdenum, from 0.5% to 1.5%manganese, from 0.1% to 0.35% vanadiurn, less than;10% carbon, from 0.8%to 1% silicon,

up to aluminum, from 1.5% to 2.8% soluble'ti- 'tanium, from 0.03% to-0.15% boron, and the balance being iron except for incidentalimpurities, the steps comprising' (1) solution treating'the alloy memberabove 1800 F. for atleasthalf anhour, (2) quenching the alloy member toa temperature below 1250" F., (3) reheating the quenched alloy member toa'first aging temperature of from 1100 'F.'to 1250 'F. and maintainingthe member at this'fi'r'st'agi'ng temperature for from 2 to 32 hours tocause nucleation of the precipitation hardening comp'onents whereby thehardness of the alloy increases from '50 to 100 D.P.H. points, ('4)slowly increasing the temperature of'the alloy member over a period offrom 2 to 32 hours until the member is at a second aging temperature inthe range of from 1300 F.'to '1450 F., and maintaining the alloy memberat thissecond aging-temperature to cause the hardness to increase at amore rapid ratethan during the first aging temperature'to a valueproportional to a value of at least 225 D.P.H. for an alloy having 1.3%soluble titanium and 330 D.P.H.'for an alloy having 2.4% solubletitanium.

0.5% boron, from 1.3 to 3.5% soluble titanium, the balance being ironexcept for incidental impurities, the steps comprising (1) solutiontreating the wrought member from M: to 8 hours at a temperature of from1700 F. to 2200 F., (2) quenching the wrought member to a temperaturebelow 1250 F., (3) reheating the wrought member to a first agingtemperature of from 1100 F. to 1250 F. and maintaining the member inthis temperature range for a period of time of from 2 to hours to'causenucleation of the precipitation hardening 20 components, (4) slowlyincreasing the temperature of the wrought member to a second temperaturefrom 1300 F. to 1450 F. over a period of from 2 to 16 hours, and 5)maintaining the wrought alloy members in this second temperature for aperiod of from 2 to 20 hours,'the

total aging period of steps (3) to (5) being from 10to hours, to providea highly hardened member.

References Cited in the file of this patent UNITED STATES PATENTSPilling et al. July 21, 1936

1. IN THE PROCESS OF PRODUCING A PRECIPITATION HARDENED MEMBER FROM ANIRON BASE AUSTENTIC ALLOY CONTAINING FROM 1.30% TO 3.5% BY WEIGHT OFSOLUBLE TITANIUM AS A PRECIPITATION HARDENING COMPONENT, THE STEPSCOMPRISING (1) SOLUTION TREATING THE ALLOY MEMBER AT A TEMPERATURE ABOVE1700*F., (2) RAPIDLY COOLING THE ALLOY MEMBER TO A TEMPERATURE BELOW1250*F., (3) REHEATING THE COOLED MEMBER TO A FIRST AGING TEMPERATURE OFFROM 1100*F. TO 1250*F. AND PREFERABLY BETWEEN 1100*F. AND 1200*F. FORALLOYS HAVING LESS THAN 1.4% OF SOLUBLE TITANIUM, AND MAINTAINING THEMEMBER IN THIS TEMPERAATURE RANGE FOR A PERIOD OF TIME OF FROM 2 TO 50HOURS SUFFICIENT TO CAUSE NUCLEATION OF THE PRECIPITATION HARDENINGCOMPONENTS WHEREBY THE HARDNESS OF THE ALLOY INCREASES FROM 30 TO 125POINTS BY THE DIAMOND PYRAMID HARDNESS TEST, (4) THEN HEATING THE ALLOYDMEMBER OVER A PERIOD OF FROM 2 TO 32 HOURS TO INCREASE ITS TEMPERATURESLOWLY TO A TEMPERATURE OF FROM ABOUAT 1300*F. TO 1450*F., ANDPREFERABLY FROM 1.4% OF SOLUBLE TITANIUM, FOR ALLOYS HAVING LESS THAN1.4% OF SOLUBLE TITANIUM AND MAINTAINING THE ALLOY MEMBER ON THIS RANGEUNTIL IT INCREASE IN HARDNESS AT LEAST 25 POINTS MORE THAN THE HARDNESSAFTER THE FIRST AGING TEMPERATURE.